Oxidation kinetics and non-Marcusian charge transfer in dimensionally confined semiconductors

Electrochemical reactions represent essential processes in fundamental chemistry that foster a wide range of applications. Although most electrochemical reactions in bulk substances can be well described by the classical Marcus-Gerischer charge transfer theory, the realistic reaction character and mechanism in dimensionally confined systems remain unknown. Here, we report the multiparametric survey on the kinetics of lateral photooxidation in structurally identical WS2 and MoS2 monolayers, where electrochemical oxidation occurs at the atomically thin monolayer edges. The oxidation rate is correlated quantitatively with various crystallographic and environmental parameters, including the density of reactive sites, humidity, temperature, and illumination fluence. In particular, we observe distinctive reaction barriers of 1.4 and 0.9 eV for the two structurally identical semiconductors and uncover an unusual non-Marcusian charge transfer mechanism in these dimensionally confined monolayers due to the limit in reactant supplies. A scenario of band bending is proposed to explain the discrepancy in reaction barriers. These results add important knowledge into the fundamental electrochemical reaction theory in low-dimensional systems.


General challenges for quantifying oxidation behavior
The grand challenge lies in the acquisition of excellent material platforms with welldefined parameter [V]. It is well known that atomic defects within a material represent a crucial crystallographic factor affecting its reaction reactivity. However, it is extremely challenging to obtain materials with well-defined [V] levels because of their non-uniform spatial distribution.
Conceptually speaking, any lattice defects are essential non-uniformity with respect to the perfect atoms staying in their lattice coordinates. In bulk materials, it is challenging or almost impossible to control the interior lattice defects in a "uniform" manner. Hence, it is impossible in bulk materials to carry out such kind of quantitative correlation of oxidation rate with [V] level, as did in 2D materials. In a word, 2D materials represent a unique platform to perform such an unprecedented study for fundamental reaction chemistry. In this sense, this work represents a really new effort for this fundamental issue.
On the other hand, although the atomic thickness and full expose of lattice atoms represents a structural advantage to control [V] levels on-demand, owing to the nature of randomness in vacancy creation, which follows the Poisson probability distribution, the distributions of [V] tend to be unconventionally broad (see Fig. 2k for the statistical data on [V] levels over 25 samples for each conditions). The intrinsic Poisson probability in spatial distribution (δ[V] ~ 20-30% in most cases) constitutes one of the primary origins of uncertainty in our experiment.
Moreover, during material degradation, defects tend to grow or spread along the sites or directions of raw defects, resulting in the behavior of clustering or aggregate of defects. Such a tendency of clustering represents another challenge to obtain data with low uncertainties, as required in most conventional experiments.

Accumulation of experimental uncertainty
As mentioned, the oxidation rate is a function of at least 4 fluctuated parameters (i.e., [V], RH, T, and F). The overall error in oxidation rate (r) is the result of the superposition of every parameter, that is, δr = δ[V] + δRH + δT + δF. Given the fact that δ[V] ~20-30% ( Fig. 2l), δRH ~5% (due to long-time drift), and δT ~5% (due to long-time drift) the accumulated error of δr may amount to 30-40% in most cases.
We have managed to minimize the experimental uncertainties by increasing the sampling numbers under each condition. In total, it took three years to prepare ca. 400 samples for this systematic/quantitative work covering 4 reaction parameters. To the best of our knowledge, there is no relevant literature available regarding the quantitative survey of oxidation rate in either 2D or conventional bulk materials.

Reliability of oxidation length extracted from PL imaging
The optical resolution of the PL imaging is about 200 nm when working under the 455 nm excitation. One may have doubts on its reliability and the accuracy of the experimental data. Hence, it would be wise to perform a complementary measurement for the estimation of oxidation length with other facilities such as line profile from AFM.
We clarify that the accuracy of PL imaging is comparable to AFM profiling in extracting the oxidation length, because it is a differential operation based on serial optical images and the relative low resolution in optics is removed during the differential operation. In Supplementary Fig. 2, we cross-checked the possible difference of length measured between PL and AFM measurements. In fact, the oxidation length determined by PL an AFM is 1.79±0.20 and 1.80±0.32 μm, respectively. Given the fluctuation of oxidation edges, the oxidation length extracted by these two facilities is quite comparable, which proves the reliability of our PL strategy.

Activation of basal oxidation mode
Previous theoretical calculations revealed that the lattice atoms on the basal planes are more stable than those at the edges. Thus, the basal oxidation mode is normally deactivated.
We found that the basal oxidation mode can be reproducibly activated in WS2 ( Supplementary Fig. 3 A remarkable merit of PL imaging over reflection style is the high contrast in images of 2D materials. Due to the atomic thickness of 1L WS2, it is difficult to discern the evolution of corrosion traces from a conventional microscope by optical reflection, as can be seen in the white-light reflection images before and after partial oxidation ( Supplementary Fig.  3a,b). However, the situation is highly improved after adopting the PL imaging strategy in tracing the progressive oxidation process. The color labeling greatly enhance the image contrast, so that even small features, such as the invisible narrow cracks across flakes in reflection mode, can be clearly discerned in the PL imaging. Such high brightness contrasts and spatial resolutions is favorable for increasing the experimental accuracy in estimating the oxidation rates. Supplementary Fig. 3c,d shows two typical PL images taken before and after oxidation for a deeply defect engineered 1L WS2. By comparing them, one can find that the cracks across the flake become wide and the triangular corrosion pits emerge from the lattice basal planes, which correspond to the trivial peripheral and activated basal oxidation modes, respectively. Interestingly, almost all the well-shaped corrosion triangles are aligned with the three edges are parallel to individual crystallographic directions ( Supplementary Fig.   3e), implying the existence of oxidation anisotropy in the basal mode.

Preparation of top-view STEM samples
The top-view STEM samples were prepared with a polymer-assisted transfer method. At first, the monolayer WS2 sheets initially exfoliated on PDMS substrate were transferred to SiO2/Si substrates by using a home-made transfer facility ( Supplementary Fig. 4a). In order to improve the adhesion between WS2 sheets and SiO2/Si substrates, the samples were then annealed at 120 ℃ for 2 min. Afterwards, PMMA polymers (used as the support for the atomically thin WS2) were spin-coated on the substrates twice at 2000 rpm a-d, Contrastive reflection and PL images recorded for a large-area 1L WS 2 flake before and after basal oxidation. In d, the randomly dispersed triangular corrosion pits emerge in the basal lattice planes, which is the characteristic of the basal oxidation mode. Note that the peripheral oxidation mode is always accompanied with the basal mode, since the activation energy of the former is lower than the latter. e, Enlarged PL image for d, in which the edges of the triangular corrosion pits are explicitly drawn and extended as guide lines for eyes, reflecting the three crystallographic directions in WS 2 as well as the oxidation anisotropy. Supplementary Fig. 4b) and the PMMA/WS2/SiO2/Si stacks were placed at room temperature for 5 hours to cure the PMMA support. In Supplementary Fig. 4c, a scotch tape with open window was used as auxiliary scaffold to peel off the PMMA/WS2 bilayer by carefully sticking it to PMMA edges. Then the five-layered tape/PMMA/WS2/SiO2/Si structure was immersed into a saturated sodium hydroxide (NaOH) solution for 5 min to detach the underlying SiO2/Si substrate by etching the SiO2 layer ( Supplementary Fig. 4d).
After that, the WS2 sheets were aligned with and gently pressed onto a copper STEM grid that was placed in advance on SiO2/Si substrate ( Supplementary Fig. 4f). The scaffold of scotch tape was then removed leaving the PMMA/WS2/STEM grid/SiO2/Si stack baked at 140 ℃ for 10 min ( Supplementary Fig. 4g), in which the PMMA support was softened to enhance the adhesion of WS2 sheets to the holey carbon film of the STEM grid. Finally, the PMMA polymer was dissolved with acetone and the WS2/STEM grid structure was achieved ( Supplementary Fig. 4h).

Optical characterization of top-view STEM sample
Supplementary Fig. 5 shows reflection and micro-zone PL images of WS2 sheets supported by a STEM grid covered with holey carbon films. Supplementary Fig. 5a-c shows the reflection images from low to high magnification ratios. In Supplementary Fig.   5b, multiple WS2 flakes with irregular shapes and a rectangular strip of metal Au can be seen, where the Au strip is used as a marker to facilitate the quick location of the WS2 flakes. Supplementary Fig. 5c focuses on the 1L WS2 flakes, which are denoted by dotted lines. Supplementary Fig. 5d shows the corresponding PL image of the 1L WS2 area. The PL images are helpful to identify the real thickness of the WS2 sheets on the STEM grids.

Effect of electron bombardment on lattice vacancies
It is well known that the high-energy electron beams have strong bombardment effect on nano-materials. Such an adverse effect becomes even serious for the atomically thin The dashed yellow rectangles represent the areas to be enlarged in the next panels. The area surrounded by dashed black lines denotes the 1L WS 2 , which can be clearly seen in the PL imaging mode. d, Corresponding PL image for the area shown in c.

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6a-d shows the STEM images taken before and after 20, 40 and 80 s, respectively. Supplementary Fig. 6e summarizes the number of sulfur vacancies versus irradiation time where a monotonous trend was revealed between them. After 20-s irradiation, the number of vacancies slightly increases from 8 to 9, while it quickly increases to 14 after 40-s irradiation, amounting to a 75% inflation. Hence, the irradiation time was set below 10 sec for all imaging areas, to minimize the bombardment effect. As compared to the intrinsic vacancies present in pristine 1L WS2, such a short exposure time brings extra errors below 10%.

STEM image filtering and lattice vacancy identification
The spherical aberration-corrected scanning transmission electron microscopy (STEM) has been demonstrated to characterize the atomic vacancies in chalcogenide monolayers for its ultrahigh spatial resolution. 1,2 The structural characterization on our WS2 samples was carried out with an aberration corrected Titan ChemiSTEM (FEI, USA). Supplementary Fig. 7a shows a typical raw image acquired with the STEM where polymer residues normally show up as areas with an enhanced brightness. Only the areas uncontaminated were selected for making statistics on the vacancy density. All raw images were processed through the Wiener filtering and average background subtracting to increase the contrasts among various atoms and lattice vacancies. Supplementary Fig. 7b,c shows typical atomically resolved STEM images before and after filtering. In the ADF imaging mode, heavy atoms would exhibit bright intensity in the greyscale images. Hence, the brightest dots and their adjacent slightly less bright ones correspond to the tungsten and sulfur atoms, respectively, and the darkest areas encircled by six tungsten and sulfur atoms can be ascribed to the centers of the hexagonal W-S atomic rings. The lattice vacancies, i.e., loss of atoms, would exhibit reduced brightness as compared to those of occupied sites, as indicated by the red arrows.
Alternatively, the atomic vacancies can be discerned by comparing the contrast profile lines. Supplementary Fig. 7d,e compares the profile lines covering a few W-S atomic pairs taken from the pristine and filtered images. It can be seen that the profile line becomes much smoother after noise filtering. Hence, the sulfur vacancy can be quickly counted by searching the sites with lower intensity. In this way, the vacancy density can also be accurately estimated.  Intensity (arb.units)

Raman characterization for excluding structural phase change
It is well known that structural phase change in TMDC materials is prone to take place under uniformly intercalative (e.g., Li and K atoms) or substitutional (e.g., Re, Tc and Mn atoms) doping that leads to collective atomic displacements (i.e., highly synchronous motion of atoms). In our case of vacancy generation by H2O2 treatment, the sulfur vacancies are randomly distributed and are relatively separated from each other.
Also, the doping level is relatively low (see the electronic characterization below). Hence, there is little chance to initiate a highly synchronous displacement of collective atoms to trigger a well-defined phase change.
To exclude the possibilities of structural phase change and flake-to-flake dependence, we performed corresponding Raman characterization on WS2 samples treated under various conditions. Supplementary Fig. 16a shows the typical Raman spectrum for a 30min treated WS2 sample, in which no any sign of structural phase change is observed. The two peaks located at 357 and 418 cm -2 can be assigned to the E2g 1 and A1g modes from the 2H phase, while no remarkable modes from 1T or 1T' phase (indicated by the red and blue arrows) are detected. We also checked the point-to-point variation by performing Raman mapping over broad areas and found no evidence of phase change, as given in Supplementary Fig. 16b In Supplementary Fig. 17, we further checked the flake-to-flake variation by comparing the Raman spectra between multiple WS2 samples treated with H2O2 for different durations. A consistent observation was made. In combination of the electronic and spectral characterization, it is therefore safe to exclude the presence of structural phase change due to electronic doping or collective atomic displacement. Convincing evidence against phase change can be obtained from the preservation of the semiconductive characteristic, because the TMDCs of a 1T or 1T' phase is generally metallic. Supplementary Fig. 18a shows the transfer curves for a same WS2 transistor successively treated for different durations from 0 to 30 min. The device exhibits a semiconductive characteristic under all conditions, although the gradual degradation of performance due to the increase of vacancies. Furthermore, the realistic doping levels after H2O2 treatment were extracted (under the criterion of 1p A as the off state) and shown in Supplementary Fig. 18b, which indicates only a moderate n-doping effect (~4×10 12 cm -2 ), in contrast to the presence of a high density of sulfur vacancies (~3×10 14 cm -2 ). We deduce that the sulfur vacancies are likely passivated by oxygen or other light chemical groups that are invisible in STEM. With this, we further guess that the residual brightness intensity in the sites of V2S (Supplementary Figs. 7e and 32e) stems from those light chemical groups that are grafted to the atoms adjacent to the vacancies.

Vacancy introduction revealed by XPS spectra
We performed XPS characterization to quantitatively estimate the densities of sulfur atoms removed after defect engineering in H2O2 soaking. In total, four WS2 monolayers under different tpts from 0 to 30 min were adopted. The untreated sample (tpt = 0 s) was used as reference and extended pretreatment durations of 10-30 min were employed for other samples to check the effect of after defect engineering.

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Supplementary Fig. 19a-d shows the XPS spectra for the W and S elements and corresponding Gaussian fittings, to estimate the atomic ratios between S and W. The main doublets from W are excited from its 2f orbit and located around 33.5 and 35.7 eV, while the doublets of S are from 2p orbit and located around 163.0 and 164.2 eV. Although the absolute magnitudes of detector counts vary with sample sizes, the relative intensity ratios between elements within each sample can be used to analyze the densities of removed sulfur atoms after defect engineering. In Supplementary Fig. 19e, the extracted S/W atomic ratio is plotted versus tpt. It linearly decreases from 1.97 to 1.56 as tpt increases from 0 to 30 min, indicating that the introduction of S vacancies is roughly proportional to tpt. Accordingly, the percentage and density of S were also estimated and given in Supplementary Fig. 19f As can be seen in Supplementary Fig. 20a, the frontal reaction boundaries are not necessarily shape in edge, that is, they are not parallel to the original flake edges, which is a natural consequence of fluctuation in oxidation rates due to non-uniform distribution of lattice defects. In most cases, the histogram from one flake (e.g., Supplementary Fig. 20a) cannot be well fitted with the Poisson statistical function. However, the histogram would be in a good statistical shape when the results from more flakes (Supplementary Fig.  20b,c,d) are merged ( Supplementary Fig. 20e). Thus, the number of samples would be slightly increased if large dispersion is encountered at different locations or samples, to reduce the overall uncertainties.
In the experiment, more than 400 samples were tested in total. On average, it took one day for characterizing one sample, including mechanical exfoliation, pretreatment, collection and processing of photooxidation data. Roughly, one and a half year was taken only to collect the photooxidation data, not to mention the time for other data and analyses,

Simulation on photothermal effect
To check the photothermal effect and investigate the local temperature of the samples under the illumination of a power density of 6 W/cm 2 , a commercial software was employed to simulate the T rise during illumination. The net optical absorptivity values for monolayers and underlying PDMS substrates are set to 5% and 0, respectively. The thermal conductivities are adopted as 116.8 and 0.16 W/(K· m) for WS2 and PDMS substrate 3 , respectively. The environmental T base is set at 300 K. As shown in Supplementary Fig.  22, the steady-state T value is only 300.08 K under the 6 W/cm 2 illumination. Only a small T rise of 0.08 K is observed.
To further verify the reliability of the software, we also simulated the case 1L WS2 on SiO2/Si substrates (9×10 6 W/cm 2 ) reported in reference 4 . The simulation shows a consistent local T of 1998 K with the reference, indicating the reliability of the software.

Service reliability due to degradation
As an application of the knowledge of oxidation kinetics, the time-dependent degradation behavior of electrical performance in 1L WS2 was recorded under different humidity conditions. Three RH levels, including < 0.1 ppm (in glovebox), 30% and 70%, were employed in experiment to check the realistic effect of humidity, as well as the existence of ~46% critical threshold for distinguishing the dry and wet oxidation mechanisms. Supplementary Fig. 22 | Spatial distribution of temperature. Simulation of the 2D temperature field under two steady states of different power densities (a: 6 W/cm 2 ; b: 9×10 6 W/cm 2 ). In the 300 K environment, the local temperature is only 300.08 K for 1L WS 2 on PDMS substrate under 6 W/cm 2 illumination, but the local temperature rises to 1998 K under 9×10 6 W/cm 2 illumination (SiO 2 as substrate). h, respectively. In the case of the oxygen-and water-free surrounding in glovebox, the oxidation rate is rather low and only ~20% reduction in device current is recorded in the 9 months (6480 h) long storage duration (Supplementary Fig. 23d). Supplementary Fig. 23e summarizes the current degradation for all the three FETs, which reveals a general trend of exponential decay in current with storage time. Clearly, the degradation rate under different RH surroundings follows the sequence: 70% > 30% > glovebox (<0.1 ppm). After ~150 h storage, the residual current ratios are about 5%, 50%, 99%, respectively. The slow degradation ratio in glovebox confirms the crucial roles of oxygen and humidity in oxidation, while the 10-fold difference in residual current between the 30% and 70% RH surroundings corroborates the existence of distinct dry and wet oxidation mechanisms that are separated by the ~46% RH threshold.

Detecting activation threshold in photon energy
To further prove the activation threshold in photon energy, we designed an experiment by using different laser sources at different stages, with the designed sequence: 450, 650, and 638 nm. The inactive 650-nm laser was intentionally arranged in the middle to ensure the convincingness of this experiment. As shown in Supplementary Fig. 24, oxidation behavior is excited in WS2 at stages I (450nm, 2.76 eV) and III (638 nm, 1.94 eV), it is inactive, in the middle, at stage II (650 nm, 1.91 eV). The result implies the threshold energy ~1.91 eV.

Control experiment in glovebox
Oxygen and moisture molecules are reported the most key species engaged in the oxidation process. 5 We also performed a control experiment by illuminating pretreated WS2 sheets in an oxygen-and moisture-free glovebox (both levels < 0.1 ppm) to check their roles.
In order to strictly control the concentrations of water and oxygen absorbates from surroundings, the WS2 sheets were exfoliated and pretreated (H2O2 solution, 30 min) in situ in separated glovebox chambers. Meanwhile, the PDMS substrates used were heated in advance to desorb molecules on its surface. Supplementary Fig. 25 shows the optical images for a 1L WS2 before and after 24-hour illumination in the glovebox. No noticeable morphologic change is seen from the reflection images. This result indicates that a sole light illumination cannot trigger the oxidation reaction and confirms the crucial roles of oxygen and moisture in the oxidation process.

Reaction rate versus substrate hydrophilicity
Previous studies have demonstrated that the surface condition of substrates has a direct impact on the chemical reactivity of the materials supported. Here we checked the effect of hydrophilicity of PDMS substrates on r1L. The PDMS is intrinsically a hydrophobic polymer. We intentionally tune its hydrophilicity by varying the exposure time in different high-energy plasma surroundings and the overall surface hydrophilicity was evaluated with the contact angle of water drops on them. The contact angle can be tuned from primitively hydrophobic ~110° to hydrophilic ~15° after a high dose exposure of Ar plasma. The surface modification changes the surface condition and moisture density absorbed, which hence remarkably tune the reaction rate. As shown in Supplementary Fig. 26, r1L increases monotonously with the contact angle exhibited by the surface-modified PDMS substrates.

Selection of humidity forms
When formulating r1L, we carefully selected the form of the parameter humidity (absolute or relative values, AH or RH) by comparing the plots of r1L versus AH and RH. Although adopting AH as the parameter to formulate r1L is physically more meaningful than RH, we found that adopting RH would result in a concise algebraic form accounting for all the T values from 20 to 26˚C. When adopting RH as the humidity parameter, all the curves converge at a critical threshold point around RH ~ 46% (Figs. 2n and 4d). Hence, the parameter RH is naturally adopted when rationalizing the quantitative relationship between r1L and humidity.

UPS characterization
To provide direct evidence, the ultraviolet photoelectron spectroscopy (UPS) technique was used to determine the realistic energy levels for 1L WS2 and MoS2, as shown in Supplementary Figs. 27 and 28. In Supplementary Fig. 27, we first verify the reliability of UPS for the Au calibrator and monolayer semiconductors. To this end, we prepared asgrown and p-doped 1L CVD WS2 and MoS2. A 40 nm gold (Au) film is electrically connected to all the disulfide samples and steel stage (Supplementary Fig. 27a)  We also verify the applicability of UPS on the as-grown and p-doped 1L CVD MoS2. The p-doping was fulfilled by evaporating a trace of an electron-drawing small molecule bis(trifluoromethylsulphonyl)imide, TSFI, on the MoS2 surfaces, to mimic the effect of Supplementary Fig. 27 | Check on accuracy and sensitivity of UPS analysis in determining the energy levels of a metal and a monolayer semiconductor. a, Diagram and real optical image for samples mounted on the steel stage for UPS measurement. b, UPS spectrum of Au calibrator. Inset: Enlarged regime around the Fermi edge. The work function of Au can be extracted from the difference between the excitation energy (hν = 21.12 eV) and the spectral width, W (calculated from cut-off region to Fermi edge). c, Comparison of the UPS spectra for as grown and p-doped 1L MoS 2 . A small E F shift around 0.18 eV is observed after doping. All spectra were collected under the electrostatic charge compensation mode to maximize the signal intensity. absorbed aqueous oxygen (which is unstable in the UPS vacuum chamber). Supplementary  Fig. 27c compares the values of energy levels for the 1L CVD MoS2 before and after pdoping. For the as-grown sample, the conduction band (EV) is located 6.01±0.04 eV below the vacuum energy (Evac), which is located just slightly higher than the HSE06 calculated value 6 (6.27 eV), indicating the reliability of the HSE06 method. From the Fermi edge, we deduce that EF-EV = 1.03±0.04 eV, indicating the neutral or slightly n-doping nature of the as-grown samples after considering the ~2 eV bandgap. In contrast, the Fermi level goes to the level at 0.81±0.04 eV above EV, featuring ca. 0.2 eV lowering after p-doping. The sensitivity in detecting the shift of EF by p-doping further confirms the high reliability of UPS for characterizing the energy levels of monolayer semiconductors.
Furthermore, we checked the sample uniformity by collecting the UPS spectra from three different areas in the doped 1L WS2 and MoS2. As shown in Supplementary Fig. 28  in EF positions between the two. The UPS data support the band bending picture that accounts for the difference in the additional energy barrier between WS2 and MoS2 during charge transfer. Supplementary

Reaction paths and band diagram
By comparing the magnitudes of simulated a values with experiment, we can first rule out two reaction paths pertinent to defect-free lattices. As shown in Fig. 3b, the barriers for the dissociation of oxygen species amount to 2.89 and 1.66 eV in the dry and wet conditions, respectively, where the oxidants originate from the molecular (O2) and anionic (O2 -) oxygen. The simulated energies are larger than the experimental values, being 2.5 eV and 1.4 eV for the dry and wet conditions, respectively (inset of Fig. 2p).
We then shift to two other reaction paths related to defective lattices where sulfur vacancies dominate (Fig. 3c). In the dry condition, molecular oxygen serves as the main oxidant and there are two barriers of 0.87 and 2.3 eV in the energy landscape. In contrast, the reaction barrier is reduced to 0.92 eV in the wet condition, where the highly active O2ions act as oxidants. The lowered activation energies are more reasonable to explain the experimentally extracted a values. Evidently, the magnitude of a is primarily determined by the exact type of oxidation agents generated under different humidity conditions. Supplementary Fig. 29 shows the formation of realistic band diagram in the electrochemical reaction between 1L WS2 and absorbed aqueous oxygen. Before humidity condensation onto WS2 (i.e., under dry surroundings), all the relevant energy levels are flat without any bending, as shown in Supplementary Fig. 29a. According to the UPS results, the positions of Ev ~ 5.8 eV and EF ~ 4.95 eV below the vacuum energy level (Evac) after considering the ~ 2 eV energy gap. Since the redox energy of (EF,redox) is 5.3 eV below Evac, a potential difference around 0.35 eV occurs between WS2 and aqueous oxygen, which leads to a upward band bending of 0.35 eV and an extra excitation barrier for charge transfer at the interface. 34 / 44

Defect engineering via an alternative oxidant
Besides the H2O2 solution, we also tested the feasibility of using other oxidants for defect engineering. We tried a Lewis acid, bis(trifluoromethane)sulfonimide (TFSI), as the oxidant for quick pretreatment. This material is known to be a strong p-dopant on 2D materials, 7 and is also proven a medium for PL enhancement.= 8,8 In this trial, the 1L WS2 was mechanically exfoliated on sapphire substrates. The TFSI powders were dissolved in nitromethane to make a 0.2 mg/ml solution. The TFSI solution was then drop-cast onto the sample and kept for 5 min at room temperature. Supplementary Fig. 30a shows the PL traced oxidation process. Similar to the samples pretreated by H2O2, this sample also exhibits peripheral oxidation behavior with the reaction fronts propagating from the edges to the central area. The PL signals are remarkably enhanced around the reaction fronts, implying a strong p-doping effect and modulation on concentration of neutral excitons in WS2 by TFSI. We intentionally terminated the oxidation and trace midway at 33 min, so that a part of unoxidized area was preserved. Then, we employed AFM to check the whole sample region, including both the oxidized and unoxidized areas. Supplementary Fig. 30b shows the AFM image covering

PL traced photooxidation for MoS2
Supplementary Fig. 31a shows contrastive white-light reflection and serial PL images boundaries spreading inwards. We note that the sharp PL boundaries allow direct discrimination between the local pristine and oxidized regions for quantitative analyses, contrasting with previous Raman and AFM methods, in which the detected signals, i.e., the variations in Raman intensity or morphology, cannot distinguish between full and partial oxidation states.    Fig. 31b plots the oxidation length, L1L, as a function of illumination time, t. In 24 min, L1L reaches up to 6 μm and leads to a roughly constant etching rate, r1L, of 0.4 ± 0.1 μm/min (inset of Supplementary Fig. 31b). To check the adhesion ability of the oxidization products, we also checked the surface morphologies before and after oxidization on a 3×5 μm 2 local area covering three regions: the pristine, oxidized, and PDMS substrate, as illustrated in Supplementary Fig. 31c,d. As can be seen, the small protrusions, with lateral sizes ranging from 20 to 100 nm, can be completely stripped after appropriate rinse in water, implying again their weak adhesive force to substrates. Such a clean strip of reaction residues constitutes a favorable merit for constructing clean and high-performance devices.
Analogous to WS2, the oxidization products are complex oxysulfides and oxides showing reduced amount of S and highly valent Mo. By comparing the XPS spectra before and after oxidation, two W +6 excitation doublets emerge at 233 and 236 eV ( Supplementary  Fig. 31e). In addition, the intensities of S doublets are sharply reduced at 162.5 and 163.8 eV, as accompanied with a new mode around 169 eV. The behavior in XPS spectra is quite similar between MoS2 and WS2.

Statistics on [V] versus tpt in MoS2
Supplementary Figs. 33−37 shows typical atomically resolved STEM images for pristine and defect engineered MoS2 samples for various pretreatment conditions. Supplementary   Fig. 38 summarizes the raw statistical [V] distribution for different tpt values. For each tpt condition, we collected images from 25 independent 4×4 nm 2 areas in statistics. The values of average [V], A, and standard deviation, D, of the statistical data were obtained with Gaussian fittings. It was found that [V] increases linearly from 2.5 to 7.1×10 13 cm -2 as tpt increases from 0 to 30 min. The generation rate is estimated to be about 1.5×10 12 cm -2 min -1 for 1L MoS2 in 30% H2O2 solution, which is 5-fold lower than that of 1L WS2, implying a slightly higher barrier for reaction between MoS2 and H2O2.